Review Article Advances in Multiferroic Nanomaterials ...downloads.hindawi.com/journals/jnm/2015/101528.pdf · ing block of nanomaterials, clusters become a candidate to assemble

Review ArticleAdvances in Multiferroic NanomaterialsAssembled with Clusters

Shifeng Zhao

School of Physical Science andTechnology and InnerMongolia Key Lab ofNanoscience andNanotechnology, InnerMongoliaUniversity,Hohhot 010021, China

Correspondence should be addressed to Shifeng Zhao; [email protected]

Received 17 June 2014; Accepted 28 October 2014

Academic Editor: Daniela Predoi

Copyright © 2015 Shifeng Zhao.This is an open access article distributed under the Creative Commons Attribution License, whichpermits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

As an entirely new perspective of multifunctional materials, multiferroics have attracted a great deal of attention. With therapidly developing micro- and nano-electro-mechanical system (MEMS&NEMS), the new kinds of micro- and nanodevicesand functionalities aroused extensive research activity in the area of multiferroics. As an ideal building block to assemble thenanostructure, cluster exhibits particular physical properties related to the cluster size at nanoscale, which is efficient in controllingthe multiferroic properties for nanomaterials. This review focuses on our recent advances in multiferroic nanomaterials assembledwith clusters. In particular, the single phase multiferroic films and compound heterostructured multiferroic films assembled withclusters were introduced detailedly.This technique presents a new and efficient method to produce the nanostructuredmultiferroicmaterials for their potential application in NEMS devices.

1. Introduction

Multiferroics have attracted increasing attention due tosimultaneous coexistence of ferromagnetic, ferroelectric, orferroelastic ordering [1, 2]. Many researchers focused on themagnetoelectric effect driven by the prospect of controllingpolarization bymagnetic field andmagnetization by electricalfield [3], which opens up an entirely new perspective of mag-netic/ferroelectric data storage media, spin-based devices(spintronics), magnetocapacitive devices, magnetic sensors,nonvolatile memories, random access memory, and so forth[4–8]. Since its discovery a century ago, ferroelectricityhas been linked to the ancient phenomena of magnetism.Attempts to combine the dipole and spin orders into onesystem started in the 1960s in Cr

2O3single crystal [9, 10],

and other single phase multiferroics, including boracites(Ni3B7O13I, Cr3B7O13Cl) [10], fluorides (BaMF

4, MMn, Fe,

Co, Ni) [11, 12], magnetite Fe3O4[13], (Y/Yb)MnO

3[14],

and BiFeO3[15], were identified in the following decades.

However, such a combination in these multiferroics has beenproven to be unexpectedly tough.

Moreover, by growing composite films combined withpiezoelectric andmagnetostrictivematerials, the strongmag-netoelectric coupling effect could be achieved due to productproperty. Much work has been done to prepare the com-posite films by combining perovskite ferroelectric oxides(e.g., Pb(Zr

0.52Ti0.48

)O3(PZT), BaTiO

3) with ferromagnetic

oxides (e.g., CoFe2O4, La0.67

Sr0.33

MnO3) [16–21]; however,

due to the low magnetostriction of these ferromagneticoxides and Ni metal, the reported magnetoelectric effects inthese composite films are generally not strong.

As well known, rare earth iron alloy (R-Fe, R = rare earthelement) possesses giant magnetostriction, being an order ofmagnitude greater than the ferromagnetic oxides [22]. Theprevious investigations have shown that the magnetoelectriceffect in the bulk laminate consisted of R-Fe alloy and fer-roelectric oxide (e.g., Tb

0.30Dy0.70

Fe2(Terfenol-D)/PZT) with

magnetoelectric coupling coefficient 𝛼𝐸of ∼4680mV/cm⋅Oe

is much larger than that of the all-oxide laminates (e.g.,CoFe2O4/PZT; its 𝛼

𝐸is ∼60mV/cm⋅Oe) [23–25]. However,

with the applications in themicro-electro-mechanical system(MEMS) devices such as microtransducers, microactuators,

Hindawi Publishing CorporationJournal of NanomaterialsVolume 2015, Article ID 101528, 12 pageshttp://dx.doi.org/10.1155/2015/101528

2 Journal of Nanomaterials

Clusters

Cluster source

SkimmerSubstrate holder

Figure 1: The sketched diagram of the growth and deposition process of clusters.

and microsensors [26], the well-defined microstructurestogether with the tunable properties are necessary for themultiferroics. Therefore, for the laminated composite film(i.e., thin-film heterostructure), it could be expected thatmagnetoelectric effect would be enhanced significantly if theR-Fe alloy is used in the magnetostriction layer. Thoughthe giant magnetostrictive films have been prepared bysong conventional film preparation means [27–30], since thephase-formation temperature of R-Fe alloy is very high (thesubstrate is generally heated above 500∘C), it is unavoidableto bring about serious oxygen diffusion from PZT oxideto Tb-Fe alloy. As a result, both magnetostriction in Tb-Fealloy and piezoelectricity in PZT are seriously suppressed.Moreover, the serious oxygen diffusion would also generatea new interface layer, which further significantly decreasesthe magnetoelectric coupling efficiency [31]. Therefore, aprogress with low temperature and energy is necessary dur-ing deposition. Fortunately, we have developed an effectivepreparation method, namely, low energy cluster beam depo-sition (LECBD), to prepare the nanostructured magnetic,giant magnetostrictive, single phase multiferroic, and well-defined microstructured multiferroic heterostructured films[32–38]. This paper aims to review the breakthroughs on themultiferroic nanostructure assembled with clusters.

2. Low Energy Cluster Beam Deposition

As a very important building block of nanomaterials, nan-oclusters are aggregates of atoms or molecules of nanometricsize, containing a number of constituent particles rangingfrom 10 to 106 [39–41]. All the beam experiments on clustershave a cluster source, in which the clusters are produced.There are a variety of sources available including arc clusterion source [42], laser vaporization cluster source [43, 44], gasaggregation source [45, 46], seeded supersonic nozzle source[47], ion sputtering source [48, 49], and liquid metal ionsource [50, 51]. Based on these we developed low energy clus-ter beam deposition method. A magnetron-sputtering-gas-aggregation (MSGA) cluster source produce is used to pro-duce cluster beam to assemble multiferroic films.The growthand deposition process of clusters is shown in Figure 1 usinggas aggregation. A direct current (DC) pulse power was usedas the sputtering power. As the sputtering gas, one stream ofargon gas with 99.9% purity was introduced through a ring

structure close to the surface of the target. Another streamof argon was fed as a buffer gas through a gas inlet nearthe magnetron discharge head.The cluster condensation andgrowth region was cooled with liquid nitrogen. A highlyoriented cluster beam with a small divergent angle less thanone degree was formed by differential pumping controlledby the skimmers. During the process of the deposition, theaverage velocity perpendicular to the substrate of each clusteris quite low (with less than 50meV/atom corresponding tokinetic energy), and the velocity parallel to the substrate isabout several meV/atom, which are both much lower thanthe molecular binding energy. The clusters are depositedon the substrate by a soft-landing manner and accumulatedrandomly but they do not coalesce with each other.

Based on this technique, various cluster-assembled nano-structured films such as metal and oxide have been prepared,which show peculiar properties different from the filmsprepared by the common methods [52–55]. Since the size,mass, and the assembling manner of the clusters can beprecisely tuned by changing the working gas flow, controllingthe length of condensation region, changing the buffer gas,and so forth, it is possible to control the microstructuresand properties of the cluster-assembled nanostructured films,whichmakes it an ideal candidate for the fabrication of singlephase or heterostructured films.

3. Magnetic Films Assembled with Clusters

3.1. Giant Magnetostrictive R-Fe Films. The cubic Lavesphase R-Fe (R = rare earth element such as Tb, Dy, Sm)compounds are well known to be a giant magnetostrictivematerial at room temperature, which could be widely usedas actuators, transducers, dampers, and so forth [56, 57].With the demand of the rapidly developing nano-electro-mechanical system (NEMS), much work has been done andvariousmethods such as ion plating [27], ion beam sputtering[28], flash evaporation [29] and magnetron sputtering [58],and molecular beam epitaxy [59, 60] have been developedfor the preparation of R-Fe films. However, compared withthe bulk materials, the saturation magnetostriction of thecurrent R-Fe films is much lower while the magnetic drivingfield is still higher [58], which limits them to be furtherused. Fortunately, based on the LECBD technique, a well-defined Tb-Fe nanostructured film has been obtained, which

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Figure 2: SEM images of the typical as-deposited Tb-Fe films and the graph of population versus size distribution of the nanoparticles atdifferent growth region length. (a) 110mm, (b) 950mm (c), and 80mm ([36]).

exhibits excellent magnetostriction, much higher than thecommon Tb-Fe films [34]. Since the size, mass, and theassemblingmanner of the R-Fe clusters can be precisely tunedby changing the working gas flow, controlling the length ofcondensation region, changing the buffer gas, and so forth,it is possible to control the microstructures and magneticproperties of the cluster-assembled R-Fe nanostructuredfilms, which makes it an ideal candidate for the fabricationof the magnetic NEMS devices.

A DC-magnetron-sputtering-gas-aggregation (MSGA)cluster source was used to produce the Tb-Fe cluster beam,which was finally deposited on the Si(100) substrate at roomtemperature to form the nanofilm. In order to tune the clustersize, the length of the condensing growth region 𝐿 was set as80mm, 95mm, and 110mm, respectively. Figure 2 presentsthe SEM images and size distributions of the typical Tb-Fe films prepared under different length of the condensinggrowth region 𝐿. One can clearly observe two facts: (i) allfilms are assembled by the spherical nanoparticles, which aredistributed uniformly and monodispersely in the film; (ii)with the increase of the condensing growth region length,the size of the nanoparticle increases, being in the rangesof 31∼36 nm for 𝐿 = 110mm, 28∼33 nm for 𝐿 = 95mm,and 23∼28 nm for 𝐿 = 80mm, respectively. Meanwhile, wenote that the particle size distribution was almost lognormal.The formation of such nanoparticle film is attributed to

the unique LECBD preparation process. And the length ofthe cluster growth region significantly influences the size ofthe Tb-Fe clusters for the films. In fact, the growth of theclusters as well as their size distribution is mainly determinedby the cluster residence time and its distribution [61, 62].With increasing the length of the condensing growth region,the cluster residence time in the condensing growth regionincreases, and thus the collision among metal ion, Tb-Fevapor, carrier gas, and free clusters becomes more sufficient,leading to the bigger size of the clusters.

It has been confirmed that all present nanoparticle-assembled Tb-Fe films exhibit higher magnetostriction com-paring to the common nonnanostructured films prepared byother methods [63–65]. And we observe that the magne-tostrictive behavior and piezomagnetic coefficient evidentlyvary with the average size of the nanoparticle. Figure 3 givesthe magnetostriction and piezomagnetic coefficient on themagnetic field for the films with various particle sizes. Withincreasing the particle size, the saturation magnetostriction𝜆𝑠and the saturation magnetic field𝐻

𝑚change, for example,

𝜆𝑠∼ 816 × 10−6 and𝐻

𝑚∼6.0 kOe for 𝑑 = 25 nm, 𝜆

𝑠∼ 1029 ×

10−6 and 𝐻𝑚∼ 7.0 kOe for 𝑑 = 30 nm, and 𝜆

𝑠= 746 ×

10−6 and 𝐻𝑚= 5.0 kOe for 𝑑 = 35 nm. Obviously, the film

with 𝑑 = 30 nm has the highest saturation magnetostrictionand piezomagnetic coefficient. However, it is not the caseat low magnetic field. The film with 𝑑 = 35 nm possesses


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highermagnetostriction and piezomagnetic coefficient at lowmagnetic field.

It suggests that the dependence of the magnetostrictionand piezomagnetic coefficient on the particle size could beattributed to the difference of the magnetization character-istic for these films. Figure 4 presents the field dependentmagnetization at room temperature for the films with theparticle sizes of 25 nm, 30 nm, and 35 nm. It is shownthat the degree of magnetization anisotropy for the film issignificantly affected by the particles size. Both in-plane andout-of-plane saturation magnetization change, as well as thecoercivity with variation of the particle size. For the filmswithparticles size of 30 nm, the degree of magnetic anisotropy isthe maximum and the difference between in-plane and out-of-plane saturationmagnetization is maximum.The easy axisis out-of-plane.Thus particle size dependence of themagneticanisotropy for the present films should be correlative to theexchange coupling effects between the nanoparticles in thefilm [66].Thus the filmwith 𝑑 = 30 nmmay show the highestdegree ofmagnetic anisotropy because the exchange couplingdistance is twice of the domain wall width (RFe∼ 15 nm)for magnetic nanoparticles [67]. Therefore, for the film with𝑑 = 30 nm, it has higher magnetic anisotropy than the otherfilms. It needs far higher magnetic field to rotate the spin intothe applied field. Therefore, the magnetostrictive coefficientof this film is lower than the other films at a low magneticfield, but, its saturationmagnetostriction is the highest, whichcould be ascribed to the higher energy of anisotropy exchangeinteraction [68].

3.2. Enhanced Ferromagnetism of BiFeO3Films Assembled

with Clusters. BiFeO3(BFO) is one of the most outstanding

single-phase lead-free multiferroics due to its high ferroelec-tric Curie (𝑇FE ∼ 1103 K) [69] and Neel (𝑇𝑁 ∼ 643K) temper-atures [70]. However, for BFOmaterials, antiferromagnetismand a superimposed incommensurate cycloid spin structure

with a periodicity of 62 nm along the [1 1 0]ℎaxis cancel

the macroscopic magnetization at room temperature, whichrestricts its applications [71]. Some investigations show thatweak ferromagnetism is observed in some limited-dimensionmaterials such as nanowires and nanoparticles due to thepartial destruction of the spiral periodicity [72–74], whichdemonstrates a possible way to enhance ferromagnetism insingle-phase multiferroics. Thus, as a size controllable build-ing block of nanomaterials, clusters become a candidate toassemble multiferroics. Therefore, using LECBD technique,we have prepared the well-defined BiFeO

3nanostructured

films assembled with 0-characteristic-dimension clusters,and then the films were annealed at 600∘C. As we expected,the ferromagnetism of the as-prepared BiFeO

3films is

enhanced [38].Figure 5 gives the morphologies of cluster-assembled

BiFeO3nanostructured films before annealing. It can be

seen that the films are assembled with clusters, which arenearly spherical and densely packed to form the uniformlycontinuous films, whereas each individual cluster is stillclearly distinguishable.The population versus the size revealsthat the average size of the nanoparticles is ∼22 nm for as-deposited films and ∼25.5 nm for the annealed films and isattributed to the fact that the size of cluster increases derivefrom the improvement of crystallizing during the annealingprocess.

Figure 6 present the XRD patterns of the typical as-deposited and annealed nanostructured films assembled withclusters. They show that both present films are polycrys-talline and all of the observed diffraction peaks can beindexed to a perovskite structure. And the as-deposited BFOnanostructured films and other common films belongingto the rhombohedral structure with space group R-3c (161)prepared by other methods annealed BFO films transformto the coexistence of tetragonal and orthorhombic symmetrystructure as (104) and (110) diffraction peaks are not obviouslysplit, which is observed by expanding the view of the XRDpattern around 2𝜃 = 32.6∘ in the inset of Figure 6. At thesame time, the lattice constant of cluster-assembled BiFeO

3

films is 𝑎 = 5.491 A, smaller than those of the films preparedby other methods. It suggests that there exists a crystaldistortion for the cluster-assembled BFO films, giving rise toa transition from the rhombohedral structure to tetragonalone [75–77], which means the crystal structure changes froma high symmetry state to a low symmetry state compared tobulk BFO materials. Thus the crystal distortion is due to thesize effect of the clusters with the smaller characteristic size,which partially destroys the long-range cycloid spin structurewith a periodicity of 62 nm in the rhombohedral structurewith space group R-3c (161). It is such crystal distortion ofthe as-prepared films that brings about the enhancement inmagnetization.

Figure 7 shows the magnetic hysteresis loops for thecluster-assembled BFO nanostructured films measured at5 K and 300K. As can be seen, obvious ferromagnetismis observed for the cluster-assembled BFO nanostructuredfilms not only at 5 K but also at room temperature. InParticular, the saturation magnetization of the BFO films atroom temperature reaches 108 emu/cc, which is comparable

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with that of the present films with 125 emu/cc at 5 K. Moreimportantly, the large magnetization of 81 emu/cc is obtainedat a magnetic field of 3000Oe at room temperature, which isa larger response than the common films prepared by othermethods [78–80].

Such enhanced room temperature ferromagnetism isattributed to the fact the average size of BFO clusters ismuch less than the long-range cycloid order of 62 nm alongthe [1 1 0]

ℎaxis that the periodicity of the spin cycloid

is broken [81]. Antiferromagnetic materials are considered

as the combination of one sublattice with spins along onedirection and anotherwith spins along the opposite direction.If no spin canting is considered, the spins of these two sub-lattices compensate each other so that the net magnetizationinside the material would become zero [82]. However, thelong-range antiferromagnetic order is frequently interruptedat the cluster surfaces, which forms the uncompensatedsurface spins. For the cluster-assembled BFO films with theaverage size of 25.5 nm, the uncompensated surface spinsbecome very significant due to the very large surface to


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volume ratio for the clusters.The uncompensated spins at thesurface enhance the contribution to the nanoparticle’s overallmagnetization. Besides, structural distortion and changeof lattice parameter due to the size effect for the cluster-assembled nanostructured films [83] lead to the release ofthe latent magnetization locked within the cycloid. Thenthe ferromagnetism of nanostructured films is significantlyenhanced.

4. Multiferroic Film HeterostructureAssembled with Clusters

It is well known that composite films combined with piezo-electric and magnetostrictive materials can obtain strongermagnetoelectric effect than single phase materials [84] by

the magnetic-mechanical-electric coupling product interac-tion via the stressmediation. Some composite films combinedby perovskite ferroelectric oxides (e.g., Pb(Zr

0.52Ti0.48

)O3

(PZT), BaTiO3) with ferromagnetic oxides (e.g., CoFe

2O4,

La0.67

Sr0.33

MnO3) [18, 21, 85, 86] did not acquire strongmag-

netoelectric effect due to low magnetostriction of the ferro-magnetic oxides. Since we had prepared giant magnetostric-tive R-Fe films using low energy cluster beam deposition, itis possible to prepare the well-defined microstructured thin-film multiferroic heterostructure consisting of R-Fe alloyand ferroelectric oxide. And the substrate is ferroelectricoxide; the degree of the interfacial reaction or diffusionbetween Tb-Fe alloy and ferroelectric oxide would be greatlysuppressed due to the low temperature and energy duringLECBD progress. Thus well-defined microstructure of thin-film heterostructure as well as strong magnetoelectric effectwould be obtained.

4.1. Tb-Fe/PZT Thin-Film Heterostructure. Tb-Fe nanoclus-ter beam was deposited onto the surface of the PZT filmthrough the open holes of the mask by LECBD progress.After deposition, not taking off the mask, a Pt electrode layerwas deposited on the Tb-Fe dots via pulse laser deposition.Figure 8 presents the surface SEM image of the Tb-Fe layer inthe heterostructure. It shows that the Tb-Fe layer is compactlyassembled by the regular spherical nanoclusters, which aredistributed uniformly and adjacent with each other. Thestructure of the thin-filmheterostructure is sketched in Insert(a) of Figure 8. Insert (b) of Figure 8 shows the cross-sectionalSEM image of the thin-film heterostructure. One observesthat the interface between Tb-Fe and PZT layers is clear andno transition layer is observed, which is benefit from theLECBD progress. During this process, the phase formationof Tb-Fe nanoclusters (or nanoparticles) is achieved in thecondensation chamber with high temperature, while thedeposition of Tb-Fe nanocluster beam onto the substrate isachieved in another high vacuum chamber with low energy


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Figure 8: The surface SEM image of the Tb-Fe layer in the thin-film heterostructure. Insert (a) is a sketch of the heterostructure,and Insert (b) is the typical cross-sectional SEM image of the het-erostructure.

and low temperature (e.g., room temperature). Both pro-cesses are independent of each other. It is easy to understandno reaction between Tb-Fe and PZT layers.

No destroying between Tb-Fe and PZT layers gives afeasibility of well ferroelectric and ferromagnetic properties.Figure 9 gives the polarization versus electric field hysteresisloops and magnetic hysteresis loops for the Tb-Fe/PZT thin-film heterostructure. It shows that the well-defined ferro-electric loops are observed. The saturation polarization andremanent polarization for Tb-Fe/PZT thin-film heterostruc-ture have a very slight decrease compared with the purePZT film. Such slight decrease in ferroelectric properties ofthe heterostructure should be attributed to the increase ofoxygen vacancy concentration in PZT layer, which bringsabout difficulty for the mobility of domain walls in a certaindegree and further leads to the decrease in polarization [87].Insert of Figure 9(a) shows that the leakage current density in

the heterostructure is quite low, for example, only being∼1.5×10−4 A/cm2 even under the higher electric field of 30MV/m.In spite of this, we found that the leakage current density inthe heterostructure was still higher than that of the pure PZTfilm, which indicates the increase of free carrier density inPZT layer of the heterostructure [88].

Besides, the heterostructure exhibits the well-definedmagnetic hysteresis loops. It shows that both in-plane andout-of-plane coercive field are the same as only 𝐻

𝑐∼

60Oe, much lower than that of the bulk Tb-Fe alloy, whilethe in-plane and out-of-plane saturation magnetizationsare, respectively, ∼38 emu/cm3 and ∼47 emu/cm3. We noticethat the magnetization character of the heterostructure isalmost comparable to the pure nanostructured Tb-Fe filmprepared by LECBD process. Since magnetoelectric effect ina two-phase composite mainly originates from the interfacialstress transfer between the magnetostrictive and the ferro-electric phase, the high magnetostriction and ferroelectricsare beneficial to the magnetoelectric coupling. A strongmagnetoelectric coupling could be obtained in the thin-filmheterostructure.

Figure 10 plots the magnetic bias𝐻bias dependence of theinduced voltage increment |Δ𝑉ME| at a given ac magneticfield frequency 𝑓 = 1.0 kHz. It shows that thin-filmheterostructure exhibits strong magnetoelectric coupling.The calculated maximum increment of the magnetoelectricvoltage coefficient is as high as ∼140mV/cm⋅Oe, larger thanthat of the reported all-oxide ferroelectric-ferromagneticcomposite film [16–18]. So strong magnetoelectric effect inTb-Fe/PZT thin-film heterostructure is evidently beneficialfrom the unique LECBD process. Based on this process,not only could the interface reaction be availably avoidedon the maximum degree, but also both ferroelectric andmagnetostrictive properties for PZT and Tb-Fe could bemaintained well.


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Insert of Figure 10 shows the 𝐻bias dependence of pie-zomagnetic coefficient 𝑞 (= 𝛿𝜆/𝛿𝐻bias) for the pure Tb-Fe nanostructured film prepared by LECBD process. Both|Δ𝑉ME| in heterostructure and 𝑞 in Tb-Fe film have thesimilar change trend with 𝐻bias. This indicates that themagnetoelectric coupling in the heterostructure should bedominated by the magnetic-mechanical-electric transformthrough the stress-mediated transfer.

4.2. Sm-Fe/PVDF Thin-Film Heterostructure. For Si basedmagnetoelectric composited films, the couple efficiencybetween ferroelectric and ferromagnetic phases was de-pressed due to the stress clamping effect of the hard substrate.Therefore a flexible polyvinylidene fluoride (PVDF) filmmaybe used instead of the hard substrate due to its small Young’smodulus. Thus the magnetoelectric coupling between theferroelectric and ferromagnetic phases will almost be notinfluenced. Moreover, piezoelectric voltage constant (𝑔

31) of

PVDF film is an order higher than that of ordinary PZT film,which allows it to generate a bigger voltage output under asmall stress.This indicates that PVDF film is suitable to act asthe piezoelectric phase in the magnetoelectric thin-film het-erostructure.

The flexible PVDF/Sm-Fe heterostructural film was pre-pared by depositing Sm-Fe nanocluster beam onto the PVDFfilm at room temperature using LECBD technique. Thoughit is very easy to destroy the PVDF polymer substrate, it canbe avoidable during LECBD progress with quite low energyand low temperature. Figure 11 shows the cross-section SEMimage of the Sm-Fe/PVDF film. It can clearly be observedthat the interface between the PVDF film and Sm-Fe layer isclear and no evident transition layer appears, indicating thatthe PVDF film does not get destroyed during the process ofLECBD. The well-defined heterostructure makes it possibleto generate the strong magnetoelectric effect.

The Sm-Fe film exhibits strong negative magnetostrictiveeffect with a saturation value of ∼750 × 10−6 at magneticfield of ∼7.0 kOe as shown in Figure 12. Inset of Figure 12shows that the Sm-Fe/PVDF film exhibits distinct magneticanisotropy with an in-plane magnetic easy axis, whichobviously makes the magnetoelectric coupling in the Sm-Fe/PVDF film be more efficient under an in-plane magneticfield.


Figure 11: The cross-section SEM image of the Sm-Fe/PVDF thin-film heterostructure.

𝜆(×10

−6)

300

0

−300

−600

M(e

mu/

cc)

60

30

0

−30

−60

H (kOe)

Out-of-plane

In-plane

H (kOe)

−16 −8 0 8 16

0 2 4 6 8

Figure 12: Magnetostriction 𝜆 dependence of Sm-Fe film onmagnetic field𝐻. Inset is the magnetic hysteresis loops for the Sm-Fe/PVDF thin-film heterostructure measured at room temperature.

For thin-film heterostructure, a large magnetoelectricvoltage output can be obtained. Figure 13 gives the magneto-electric voltage output increment |Δ𝑉ME| value as a functionof 𝐻bias for the Sm-Fe/PVDF film. It is seen that the filmexhibits a large voltage output under the external magneticbias. The |Δ𝑉ME| value increases with increasing𝐻bias, reach-ing themaximumvalue of |Δ𝑉ME| ∼ 210 𝜇Vat𝐻bias = 2.3 kOe,and then drops. Compared with the previous investiga-tions, the magnetoelectric voltage output in the present Sm-Fe/PVDF film is remarkably large, almost being two ordershigher than that of typical all-oxide PZT/CoFe

2O4/PZT film

deposited on the hard wafer [89].Therefore, by using the flexible PVDF polymer film as

the substrate, the substrate clamping effect on the magne-toelectric coupling of the heterostructural film is completelyeliminated. The heterostructural film exhibits large magne-toelectric voltage output, which is mainly attributed to thelarge piezoelectric voltage constant in the piezoelectric PVDFlayer and high magnetic anisotropy with in-plane magneticeasy axis as well as the giant negative magnetostriction in theferromagnetic Sm-Fe layer.

200

100

0

0 2 4 6

Hbias (kOe)

|ΔV

ME|

(𝜇V

)

Figure 13: Induced voltage increment |Δ𝑉ME| as a function of mag-netic bias𝐻bias.

5. Conclusions

In a conclusion, the single phase multiferroic films and com-pound heterostructured multiferroic films assembled withclusterswere prepared by low energy cluster beamdeposition.It shows that the multiferroic properties of the thin-films canbe controlled or improved by tuning the size of the clusters.And the structure of the thin-film heterostructure wouldnot be destroyed due to low temperature and energy duringLECBD progress. The LECBD technique provides an idealavenue to prepare multiferroic nanostructure and facilitatestheir applications on NEMS devices.

Conflict of Interests

The author declares that there is no conflict of interestsregarding the publication of this paper.

Acknowledgments

This work was financially supported by National Key Projectsfor Basic Research of China (973 Projects) (Grant no.2012CB626815), the National Natural Science Foundation ofChina (Grant nos. 11264026, 10904065), the Program forYoung Talents of Science and Technology in Universities ofInner Mongolia Autonomous Region (Grant no. NJYT-12-B05), and Inner Mongolia Science Foundation for Distin-guished Young Scholars (Grant no. 2014JQ01).

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